Ceracon forging of l12 aluminum alloys

ABSTRACT

A method for producing high strength aluminum alloy consolidated billets containing L1 2  dispersoids by Ceracon forging is disclosed. The method comprises forming an aluminum alloy powder compact preform containing L1 2  dispersoid forming elements therein and encompassing the preform in a flowable pressure transmitting medium in a die in a hydraulic press. The die, pressure transmitting medium and preform are then heated and the preform is forged by applying pressure to the pressure transmitting medium by the ram of the hydraulic press. The unequal axial and radial strain resulting from this type of forging results in improved mechanical properties of L1 2  aluminum alloys.

CROSS-REFERENCE TO RELATED APPLICATION(S)

This application is related to the following co-pending applicationsthat were filed on Dec. 9, 2008 herewith and are assigned to the sameassignee: CONVERSION PROCESS FOR HEAT TREATABLE L1₂ ALUMINUM ALLOYS,Ser. No. 12/316,020; A METHOD FOR FORMING HIGH STRENGTH ALUMINUM ALLOYSCONTAINING L1₂ INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,046; and AMETHOD FOR PRODUCING HIGH STRENGTH ALUMINUM ALLOY POWDER CONTAINING L1₂INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,047.

This application is also related to the following co-pendingapplications that were filed on Apr. 18, 2008, and are assigned to thesame assignee: L1₂ ALUMINUM ALLOYS WITH BIMODAL AND TRIMODALDISTRIBUTION, Ser. No. 12/148,395; DISPERSION STRENGTHENED L1₂ ALUMINUMALLOYS, Ser. No. 12/148,432; HEAT TREATABLE L1₂ ALUMINUM ALLOYS, Ser.No. 12/148,383; HIGH STRENGTH L1₂ ALUMINUM ALLOYS, Ser. No. 12/148,394;HIGH STRENGTH L1₂ ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLEL1₂ ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH L1₂ ALUMINUMALLOYS, Ser. No. 12/148,387; HIGH STRENGTH ALUMINUM ALLOYS WITH L1₂PRECIPITATES, Ser. No. 12/148,426; HIGH STRENGTH L1₂ ALUMINUM ALLOYS,Ser. No. 12/148,459; and L1₂ STRENGTHENED AMORPHOUS ALUMINUM ALLOYS,Ser. No. 12/148,458.

BACKGROUND

The present invention relates generally to aluminum alloys and morespecifically to a method for forming high strength aluminum alloybillets having L1₂ dispersoids therein.

The combination of high strength, ductility, and fracture toughness, aswell as low density, make aluminum alloys natural candidates foraerospace and space applications. However, their use is typicallylimited to temperatures below about 300° F. (149° C.) since mostaluminum alloys start to lose strength in that temperature range as aresult of coarsening of strengthening precipitates.

The development of aluminum alloys with improved elevated temperaturemechanical properties is a continuing process. Some attempts haveincluded aluminum-iron and aluminum-chromium based alloys such asAl—Fe—Ce, Al—Fe—V—Si, Al—Fe—Ce—W, and Al—Cr—Zr—Mn that containincoherent dispersoids. These alloys, however, also lose strength atelevated temperatures due to particle coarsening. In addition, thesealloys exhibit ductility and fracture toughness values lower than othercommercially available aluminum alloys.

Other attempts have included the development of mechanically alloyedAl—Mg and Al—Ti alloys containing ceramic dispersoids. These alloysexhibit improved high temperature strength due to the particledispersion, but the ductility and fracture toughness are not improved.

U.S. Pat. No. 6,248,453 owned by the assignee of the present inventiondiscloses aluminum alloys strengthened by dispersed Al₃X L1₂intermetallic phases where X is selected from the group consisting ofSc, Er, Lu, Yb, Tm, and Lu. The Al₃X particles are coherent with thealuminum alloy matrix and are resistant to coarsening at elevatedtemperatures. The improved mechanical properties of the discloseddispersion strengthened L1₂ aluminum alloys are stable up to 572° F.(300° C.). U.S. Patent Application Publication No. 2006/0269437 Al alsoowned commonly discloses a high strength aluminum alloy that containsscandium and other elements that is strengthened by L1₂ dispersoids.

L1₂ strengthened aluminum alloys have high strength and improved fatigueproperties compared to commercially available aluminum alloys. Finegrain size results in improved mechanical properties of materials.Hall-Petch strengthening has been known for decades where strengthincreases as grain size decreases. An optimum grain size for optimumstrength is in the nano range of about 30 to 100 nm. These alloys alsohave lower ductility.

SUMMARY

The present invention is a method for forming aluminum alloy billetswith high strength and acceptable fracture toughness. In embodiments,the alloys have coherent L1₂ Al₃X dispersoids where X is at least onefirst element selected from scandium, erbium, thulium, ytterbium, andlutetium, and at least one second element selected from gadolinium,yttrium, zirconium, titanium, hafnium, and niobium. The balance issubstantially aluminum containing at least one alloying element selectedfrom silicon, magnesium, lithium, copper, zinc, and nickel.

The alloys are formed by encompassing a powder preform of an aluminumalloy body containing L1₂ dispersoid forming elements in a heated,flowable pressure transmitting medium, and rapidly compressing thepowder to consolidate the perform to form the billet. Use of graphite ora ceramic as the pressure transfer medium causes a non-isostaticpressure field to form in the chamber. During consolidation, the powderpreform undergoes an axial compression that exceeds radial expansion.The resulting shear strain cleans the surface oxide from the particlesand increases metal-to-metal contact improving forging density.

This method is known in the industry as Ceracon forging. It has not beenattempted using aluminum alloy containing L1₂ dispersoids. Examples ofthe Ceracon forging process are shown in U.S. Pat. No. 4,667,497, U.S.Patent Application Publication No. 2005/0147520 and U.S. Pat. No.7,097,807 and are included in their entirety by reference.

The compression takes place in a closed container and results in billetsof aluminum alloy containing L1₂ dispersoids with a density ofessentially 100%.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is an aluminum scandium phase diagram.

FIG. 2 is an aluminum erbium phase diagram.

FIG. 3 is an aluminum thulium phase diagram.

FIG. 4 is an aluminum ytterbium phase diagram.

FIG. 5 is an aluminum lutetium phase diagram.

FIG. 6 is a schematic diagram of a vertical gas atomizer.

FIG. 7 is a scanning electron micrograph of the gas atomized inventiveL1₂ aluminum alloy powder.

FIG. 8 is a diagram showing the processing steps to consolidate L1₂aluminum alloy powder.

FIG. 9 is a schematic diagram illustrating non-isostatic forging.

FIG. 10 is a diagram showing the processing steps to Ceracon forge anL1₂ aluminum alloy powder preform.

DETAILED DESCRIPTION 1. L1₂ Aluminum Alloys

The alloy products of this invention are formed from aluminum basedalloys with high strength and fracture toughness for applications attemperatures from about −420° F. (−251 ° C.) up to about 650° F. (343°C.). The aluminum alloy comprises a solid solution of aluminum and atleast one element selected from silicon, magnesium, lithium, copper,zinc, and nickel strengthened by L1₂ coherent precipitates where X is atleast one first element selected from scandium, erbium, thulium,ytterbium, and lutetium, and at least one second element selected fromgadolinium, yttrium, zirconium, titanium, hafnium, and niobium.

The aluminum silicon system is a simple eutectic alloy system with aeutectic reaction at 12.5 weight percent silicon and 1077° F. (577° C.).There is little solubility of silicon in aluminum at temperatures up to930° F. (500° C.) and none of aluminum in silicon. However, thesolubility can be extended significantly by utilizing rapidsolidification techniques

The binary aluminum magnesium system is a simple eutectic at 36 weightpercent magnesium and 842° F. (450° C.). There is complete solubility ofmagnesium and aluminum in the rapidly solidified inventive alloysdiscussed herein

The binary aluminum lithium system is a simple eutectic at 8 weightpercent lithium and 1105° (596° C.). The equilibrium solubility of 4weight percent lithium can be extended significantly by rapidsolidification techniques. There can be complete solubility of lithiumin the rapid solidified inventive alloys discussed herein.

The binary aluminum copper system is a simple eutectic at 32 weightpercent copper and 1018° F. (548° C.). There can be complete solubilityof copper in the rapidly solidified inventive alloys discussed herein.

The aluminum zinc binary system is a eutectic alloy system involving amonotectoid reaction and a miscibility gap in the solid state. There isa eutectic reaction at 94 weight percent zinc and 718° F. (381° C.).Zinc has maximum solid solubility of 83.1 weight percent in aluminum at717.8° F. (381° C.) which can be extended by rapid solidificationprocesses. Decomposition of the super saturated solid solution of zincin aluminum gives rise to spherical and ellipsoidal GP zones which arecoherent with the matrix and act to strengthen the alloy.

The aluminum nickel binary system is a simple eutectic at 5.7 weightpercent nickel and 1183.8° F. (639.9° C.). There is little solubility ofnickel in aluminum. However, the solubility can be extendedsignificantly by utilizing rapid solidification processes. Theequilibrium phase in the aluminum nickel eutectic system is L1₂intermetallic Al₃Ni.

In the aluminum based alloys disclosed herein, scandium, erbium,thulium, ytterbium, and lutetium are potent strengtheners that have lowdiffusivity and low solubility in aluminum. All these elements formequilibrium Al₃X intermetallic dispersoids where X is at least one ofscandium, erbium, thulium, ytterbium, and lutetium, that have an L1₂structure that is an ordered face centered cubic structure with the Xatoms located at the corners and aluminum atoms located on the cubefaces of the unit cell.

Scandium forms Al₃Sc dispersoids that are fine and coherent with thealuminum matrix. Lattice parameters of aluminum and Al₃Sc are very close(0.405 nm and 0.410 nm respectively), indicating that there is minimalor no driving force for causing growth of the Al₃Sc dispersoids. Thislow interfacial energy makes the Al₃Sc dispersoids thermally stable andresistant to coarsening up to temperatures as high as about 842° F.(450° C.). Additions of magnesium in aluminum increase the latticeparameter of the aluminum matrix, and decrease the lattice parametermismatch further increasing the resistance of the Al₃Sc to coarsening.Additions of zinc, copper, lithium, silicon, and nickel provide solidsolution and precipitation strengthening in the aluminum alloys. In thealloys of this invention these Al₃Sc dispersoids are made stronger andmore resistant to coarsening at elevated temperatures by adding suitablealloying elements such as gadolinium, yttrium, zirconium, titanium,hafnium, niobium, or combinations thereof, that enter Al₃Sc in solution.

Erbium forms Al₃Er dispersoids in the aluminum matrix that are fine andcoherent with the aluminum matrix. The lattice parameters of aluminumand Al₃Er are close (0.405 nm and 0.417 nm respectively), indicatingthere is minimal driving force for causing growth of the Al₃Erdispersoids. This low interfacial energy makes the Al₃Er dispersoidsthermally stable and resistant to coarsening up to temperatures as highas about 842° F. (450° C.). Additions of magnesium in aluminum increasethe lattice parameter of the aluminum matrix, and decrease the latticeparameter mismatch further increasing the resistance of the Al₃Er tocoarsening. Additions of zinc, copper, lithium, silicon, and nickelprovide solid solution and precipitation strengthening in the aluminumalloys. In the alloys of this invention, these Al₃Er dispersoids aremade stronger and more resistant to coarsening at elevated temperaturesby adding suitable alloying elements such as gadolinium, yttrium,zirconium, titanium, hafnium, niobium, or combinations thereof thatenter Al₃Er in solution.

Thulium forms metastable Al₃Tm dispersoids in the aluminum matrix thatare fine and coherent with the aluminum matrix. The lattice parametersof aluminum and Al₃Tm are close (0.405 nm and 0.420 nm respectively),indicating there is minimal driving force for causing growth of theAl₃Tm dispersoids. This low interfacial energy makes the Al₃Tmdispersoids thermally stable and resistant to coarsening up totemperatures as high as about 842° F. (450° C.). Additions of magnesiumin aluminum increase the lattice parameter of the aluminum matrix, anddecrease the lattice parameter mismatch further increasing theresistance of the Al₃Tm to coarsening. Additions of zinc, copper,lithium, silicon, and nickel provide solid solution and precipitationstrengthening in the aluminum alloys. In the alloys of this inventionthese Al₃Tm dispersoids are made stronger and more resistant tocoarsening at elevated temperatures by adding suitable alloying elementssuch as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, orcombinations thereof that enter Al₃Tm in solution.

Ytterbium forms Al₃Yb dispersoids in the aluminum matrix that are fineand coherent with the aluminum matrix. The lattice parameters of Al andAl₃Yb are close (0.405 nm and 0.420 nm respectively), indicating thereis minimal driving force for causing growth of the Al₃Yb dispersoids.This low interfacial energy makes the Al₃Yb dispersoids thermally stableand resistant to coarsening up to temperatures as high as about 842° F.(450° C.). Additions of magnesium in aluminum increase the latticeparameter of the aluminum matrix, and decrease the lattice parametermismatch further increasing the resistance of the Al₃Yb to coarsening.Additions of zinc, copper, lithium, silicon, and nickel provide solidsolution and precipitation strengthening in the aluminum alloys. In thealloys of this invention, these Al₃Yb dispersoids are made stronger andmore resistant to coarsening at elevated temperatures by adding suitablealloying elements such as gadolinium, yttrium, zirconium, titanium,hafnium, niobium, or combinations thereof that enter Al₃Yb in solution.

Lutetium forms Al₃Lu dispersoids in the aluminum matrix that are fineand coherent with the aluminum matrix. The lattice parameters of Al andAl₃Lu are close (0.405 nm and 0.419 nm respectively), indicating thereis minimal driving force for causing growth of the Al₃Lu dispersoids.This low interfacial energy makes the Al₃Lu dispersoids thermally stableand resistant to coarsening up to temperatures as high as about 842° F.(450° C.). Additions of magnesium in aluminum increase the latticeparameter of the aluminum matrix, and decrease the lattice parametermismatch further increasing the resistance of the Al₃Lu to coarsening.Additions of zinc, copper, lithium, silicon, and nickel provide solidsolution and precipitation strengthening in the aluminum alloys. In thealloys of this invention, these Al₃Lu dispersoids are made stronger andmore resistant to coarsening at elevated temperatures by adding suitablealloying elements such as gadolinium, yttrium, zirconium, titanium,hafnium, niobium, or mixtures thereof that enter Al₃Lu in solution.

Gadolinium forms metastable Al₃Gd dispersoids in the aluminum matrixthat are stable up to temperatures as high as about 842° F. (450° C.)due to their low diffusivity in aluminum. The Al₃Gd dispersoids have aD0₁₉ structure in the equilibrium condition. Despite its large atomicsize, gadolinium has fairly high solubility in the Al₃X intermetallicdispersoids (where X is scandium, erbium, thulium, ytterbium orlutetium). Gadolinium can substitute for the X atoms in Al₃Xintermetallic, thereby forming an ordered L1₂ phase which results inimproved thermal and structural stability.

Yttrium forms metastable Al₃Y dispersoids in the aluminum matrix thathave an L1₂ structure in the metastable condition and a D0₁₉ structurein the equilibrium condition. The metastable Al₃Y dispersoids have a lowdiffusion coefficient which makes them thermally stable and highlyresistant to coarsening. Yttrium has a high solubility in the Al₃Xintermetallic dispersoids allowing large amounts of yttrium tosubstitute for X in the Al₃X L1₂ dispersoids which results in improvedthermal and structural stability.

Zirconium forms Al₃Zr dispersoids in the aluminum matrix that have anL1₂ structure in the metastable condition and D0₂₃ structure in theequilibrium condition. The metastable Al₃Zr dispersoids have a lowdiffusion coefficient which makes them thermally stable and highlyresistant to coarsening. Zirconium has a high solubility in the Al₃Xdispersoids allowing large amounts of zirconium to substitute for X inthe Al₃X dispersoids, which results in improved thermal and structuralstability.

Titanium forms Al₃Ti dispersoids in the aluminum matrix that have an L1₂structure in the metastable condition and DO₂₂ structure in theequilibrium condition. The metastable Al₃Ti despersoids have a lowdiffusion coefficient which makes them thermally stable and highlyresistant to coarsening. Titanium has a high solubility in the Al₃Xdispersoids allowing large amounts of titanium to substitute for X inthe Al₃X dispersoids, which result in improved thermal and structuralstability.

Hafnium forms metastable Al₃Hf dispersoids in the aluminum matrix thathave an L1₂ structure in the metastable condition and a D0₂₃ structurein the equilibrium condition. The Al₃Hf dispersoids have a low diffusioncoefficient, which makes them thermally stable and highly resistant tocoarsening. Hafnium has a high solubility in the Al₃X dispersoidsallowing large amounts of hafnium to substitute for scandium, erbium,thulium, ytterbium, and lutetium in the above mentioned Al₃Xdispersoids, which results in stronger and more thermally stabledispersoids.

Niobium forms metastable Al₃Nb dispersoids in the aluminum matrix thathave an L1₂ structure in the metastable condition and a D0₂₂ structurein the equilibrium condition. Niobium has a lower solubility in the Al₃Xdispersoids than hafnium or yttrium, allowing relatively lower amountsof niobium than hafnium or yttrium to substitute for X in the Al₃Xdispersoids. Nonetheless, niobium can be very effective in slowing downthe coarsening kinetics of the Al₃X dispersoids because the Al₃Nbdispersoids are thermally stable. The substitution of niobium for X inthe above mentioned Al₃X dispersoids results in stronger and morethermally stable dispersoids.

Al₃X L1₂ precipitates improve elevated temperature mechanical propertiesin aluminum alloys for two reasons. First, the precipitates are orderedintermetallic compounds. As a result, when the particles are sheared byglide dislocations during deformation, the dislocations separate intotwo partial dislocations separated by an anti-phase boundary on theglide plane. The energy to create the anti-phase boundary is the originof the strengthening. Second, the cubic L1₂ crystal structure andlattice parameter of the precipitates are closely matched to thealuminum solid solution matrix. This results in a lattice coherency atthe precipitate/matrix boundary that resists coarsening. The lack of aninterphase boundary results in a low driving force for particle growthand resulting elevated temperature stability. Alloying elements in solidsolution in the dispersed strengthening particles and in the aluminummatrix that tend to decrease the lattice mismatch between the matrix andparticles will tend to increase the strengthening and elevatedtemperature stability of the alloy.

L1₂ phase strengthened aluminum alloys are important structuralmaterials because of their excellent mechanical properties and thestability of these properties at elevated temperature due to theresistance of the coherent dispersoids in the microstructure to particlecoarsening. The mechanical properties are optimized by maintaining ahigh volume fraction of L1₂ dispersoids in the microstructure. The L1₂dispersoid concentration following aging scales as the amount of L1₂phase forming elements in solid solution in the aluminum alloy followingquenching. Examples of L1₂ phase forming elements include but are notlimited to Sc, Er, Th, Yb, and Lu. The concentration of alloyingelements in solid solution in alloys cooled from the melt is directlyproportional to the cooling rate. Exemplary aluminum alloys for thebimodal system alloys of this invention include, but are not limited to(in weight percent unless otherwise specified):

about Al—M-(0.1-4)Sc-(0.1-20)Gd;

about Al—M-(0.1-20)Er-(0.1-20)Gd;

about Al—M-(0.1-15)Tm-(0.1-20)Gd;

about Al—M-(0.1-25)Yb-(0.1-20)Gd;

about Al—M-(0.1-25)Lu-(0.1-20)Gd;

about Al—M-(0.1-4)Sc-(0.1-20)Y;

about Al—M-(0.1-20)Er-(0.1-20)Y;

about Al—M-(0.1-15)Tm-(0.1-20)Y;

about Al—M-(0.1-25)Yb-(0.1-20)Y;

about Al—M-(0.1-25)Lu-(0.1-20)Y;

about Al—M-(0.1-4)Sc-(0.05-4)Zr;

about Al—M-(0.1-20)Er-(0.05-4)Zr;

about Al—M-(0.1-15)Tm-(0.05-4)Zr;

about Al—M-(0.1-25)Yb-(0.05-4)Zr;

about Al—M-(0.1-25)Lu-(0.05-4)Zr;

about Al—M-(0.1-4)Sc-(0.05-10)Ti;

about Al—M-(0.1-20)Er-(0.05-10)Ti;

about Al—M-(0.1-15)Tm-(0.05-10)Ti;

about Al—M-(0.1-25)Yb-(0.05-10)Ti;

about Al—M-(0.1-25)Lu-(0.05-10)Ti;

about Al—M-(0.1-4)Sc-(0.05-10)Hf;

about Al—M-(0.1-20)Er-(0.05-10)Hf;

about Al—M-(0.1-15)Tm-(0.05-10)Hf;

about Al—M-(0.1-25)Yb-(0.05-10)Hf;

about Al—M-(0.1-25)Lu-(0.05-10)Hf;

about Al—M-(0.1-4)Sc-(0.05-5)Nb;

about Al—M-(0.1-20)Er-(0.05-5)Nb;

about Al—M-(0.1-15)Tm-(0.05-5)Nb;

about Al—M-(0.1-25)Yb-(0.05-5)Nb; and

about Al—M-(0.1-25)Lu-(0.05-5)Nb.

M is at least one of about (4-25) weight percent silicon, (1-8) weightpercent magnesium, (0.5-3) weight percent lithium, (0.2-6.5) weightpercent copper, (3-12) weight percent zinc, and (1-12) weight percentnickel.

The amount of silicon present in the fine grain matrix of this inventionif any may vary from about 4 to about 25 weight percent, more preferablyfrom about 4 to about 18 weight percent, and even more preferably fromabout 5 to about 11 weight percent.

The amount of magnesium present in the fine grain matrix of thisinvention if any may vary from about 1 to about 8 weight percent, morepreferably from about 3 to about 7.5 weight percent, and even morepreferably from about 4 to about 6.5 weight percent.

The amount of lithium present in the fine grain matrix of this inventionif any may vary from about 0.5 to about 3 weight percent, morepreferably from about 1 to about 2.5 weight percent, and even morepreferably from about 1 to about 2 weight percent.

The amount of copper present in the fine grain matrix of this inventionif any may vary from about 0.2 to about 6.5 weight percent, morepreferably from about 0.5 to about 5.0 weight percent, and even morepreferably from about 2 to about 4.5 weight percent.

The amount of zinc present in the fine grain matrix of this invention ifany may vary from about 3 to about 12 weight percent, more preferablyfrom about 4 to about 10 weight percent, and even more preferably fromabout 5 to about 9 weight percent.

The amount of nickel present in the fine grain matrix of this inventionif any may vary from about 1 to about 12 weight percent, more preferablyfrom about 2 to about 10 weight percent, and even more preferably fromabout 4 to about 10 weight percent.

The amount of scandium present in the fine grain matrix of thisinvention if any may vary from 0.1 to about 4 weight percent, morepreferably from about 0.1 to about 3 weight percent, and even morepreferably from about 0.2 to about 2.5 weight percent. The Al—Sc phasediagram shown in FIG. 1 indicates a eutectic reaction at about 0.5weight percent scandium at about 1219° F. (659° C.) resulting in a solidsolution of scandium and aluminum and Al₃Sc dispersoids. Aluminum alloyswith less than 0.5 weight percent scandium can be quenched from the meltto retain scandium in solid solution that may precipitate as dispersedL1₂ intermetallic Al₃Sc following an aging treatment. Alloys withscandium in excess of the eutectic composition (hypereutectic alloys)can only retain scandium in solid solution by rapid solidificationprocessing (RSP) where cooling rates are in excess of about 10³°C./second.

The amount of erbium present in the fine grain matrix of this invention,if any, may vary from about 0.1 to about 20 weight percent, morepreferably from about 0.3 to about 15 weight percent, and even morepreferably from about 0.5 to about 10 weight percent. The Al—Er phasediagram shown in FIG. 2 indicates a eutectic reaction at about 6 weightpercent erbium at about 1211° F. (655° C.). Aluminum alloys with lessthan about 6 weight percent erbium can be quenched from the melt toretain erbium in solid solutions that may precipitate as dispersed L1₂intermetallic Al₃Er following an aging treatment. Alloys with erbium inexcess of the eutectic composition can only retain erbium in solidsolution by rapid solidification processing (RSP) where cooling ratesare in excess of about 10³° C./second.

The amount of thulium present in the alloys of this invention, if any,may vary from about 0.1 to about 15 weight percent, more preferably fromabout 0.2 to about 10 weight percent, and even more preferably fromabout 0.4 to about 6 weight percent. The Al—Tm phase diagram shown inFIG. 3 indicates a eutectic reaction at about 10 weight percent thuliumat about 1193° F. (645° C.). Thulium forms metastable Al₃Tm dispersoidsin the aluminum matrix that have an L1₂ structure in the equilibriumcondition. The Al₃Tm dispersoids have a low diffusion coefficient whichmakes them thermally stable and highly resistant to coarsening. Aluminumalloys with less than 10 weight percent thulium can be quenched from themelt to retain thulium in solid solution that may precipitate asdispersed metastable L1₂ intermetallic Al₃Tm following an agingtreatment. Alloys with thulium in excess of the eutectic composition canonly retain Tm in solid solution by rapid solidification processing(RSP) where cooling rates are in excess of about 10³° C./second.

The amount of ytterbium present in the alloys of this invention, if any,may vary from about 0.1 to about 25 weight percent, more preferably fromabout 0.3 to about 20 weight percent, and even more preferably fromabout 0.4 to about 10 weight percent. The Al—Yb phase diagram shown inFIG. 4 indicates a eutectic reaction at about 21 weight percentytterbium at about 1157° F. (625° C.). Aluminum alloys with less thanabout 21 weight percent ytterbium can be quenched from the melt toretain ytterbium in solid solution that may precipitate as dispersed L1₂intermetallic Al₃Yb following an aging treatment. Alloys with ytterbiumin excess of the eutectic composition can only retain ytterbium in solidsolution by rapid solidification processing (RSP) where cooling ratesare in excess of about 10³° C./second.

The amount of lutetium present in the alloys of this invention, if any,may vary from about 0.1 to about 25 weight percent, more preferably fromabout 0.3 to about 20 weight percent, and even more preferably fromabout 0.4 to about 10 weight percent. The Al—Lu phase diagram shown inFIG. 5 indicates a eutectic reaction at about 11.7 weight percent Lu atabout 1202° F. (650° C.). Aluminum alloys with less than about 11.7weight percent lutetium can be quenched from the melt to retain Lu insolid solution that may precipitate as dispersed L1₂ intermetallic Al₃Lufollowing an aging treatment. Alloys with Lu in excess of the eutecticcomposition can only retain Lu in solid solution by rapid solidificationprocessing (RSP) where cooling rates are in excess of about 10³°C./second.

The amount of gadolinium present in the alloys of this invention, ifany, may vary from about 0.1 to about 20 weight percent, more preferablyfrom about 0.3 to about 15 weight percent, and even more preferably fromabout 0.5 to about 10 weight percent.

The amount of yttrium present in the alloys of this invention, if any,may vary from about 0.1 to about 20 weight percent, more preferably fromabout 0.3 to about 15 weight percent, and even more preferably fromabout 0.5 to about 10 weight percent.

The amount of zirconium present in the alloys of this invention, if any,may vary from about 0.05 to about 4 weight percent, more preferably fromabout 0.1 to about 3 weight percent, and even more preferably from about0.3 to about 2 weight percent.

The amount of titanium present in the alloys of this invention, if any,may vary from about 0.05 to about 10 weight percent, more preferablyfrom about 0.2 to about 8 weight percent, and even more preferably fromabout 0.4 to about 4 weight percent.

The amount of hafnium present in the alloys of this invention, if any,may vary from about 0.05 to about 10 weight percent, more preferablyfrom about 0.2 to about 8 weight percent, and even more preferably fromabout 0.4 to about 5 weight percent.

The amount of niobium present in the alloys of this invention, if any,may vary from about 0.05 to about 5 weight percent, more preferably fromabout 0.1 to about 3 weight percent, and even more preferably from about0.2 to about 2 weight percent.

In order to have the best properties for the fine grain matrix of thisinvention, it is desirable to limit the amount of other elements.Specific elements that should be reduced or eliminated include no morethan about 0.1 weight percent iron, 0.1 weight percent chromium, 0.1weight percent manganese, 0.1 weight percent vanadium, and 0.1 weightpercent cobalt. The total quantity of additional elements should notexceed about 1% by weight, including the above listed impurities andother elements.

2. Ceracon Forging of L1₂ Aluminum Alloys

It is advantageous to form L1₂ strengthened aluminum alloy product frompowder. The major reason is that the rapid cooling rate experiencedduring powder formation from the melt results in high supersaturation ofintermetallic L1₂ phase forming elements in the powder. The highsupersaturation leads to a maximum amount of the strengthening phasedispersed throughout the structure in the final consolidated part.

The highest cooling rates observed in commercially viable processes areachieved by gas atomization of molten metals to produce powder. Gasatomization is a two fluid process wherein a stream of molten metal isdisintegrated by a high velocity gas stream. The end result is that theparticles of molten metal eventually become spherical due to surfacetension and finely solidify in powder form. Heat from the liquiddroplets is transferred to the atomization gas by convection. Thesolidification rates, depending on the gas and the surroundingenvironment, can be very high and can exceed 10⁶° C./second. Coolingrates greater than 10³° C./second are typically specified to ensuresupersaturation of alloying elements in gas atomized L1₂ aluminum alloypowder in the inventive process described herein.

A schematic of typical vertical gas atomizer 100 is shown in FIG. 6A.FIG. 6A is taken from R. Germain, Powder Metallurgy Science SecondEdition MPIF (1994) (chapter 3, p. 101) and is incorporated herein byreference. Vacuum or inert gas induction melter 102 is positioned at thetop of free flight chamber 104. Vacuum induction melter 102 containsmelt 106 which flows by gravity or gas overpressure through nozzle 108.A close up view of nozzle 108 is shown in FIG. 6B. Melt 106 entersnozzle 108 and flows downward till it meets high pressure gas streamfrom gas source 110 where it is transformed into a spray of droplets.The droplets eventually become spherical due to surface tension andrapidly solidify into spherical powder 112 which collects in collectionchamber 114. The gas recirculates through cyclone collector 116 whichcollects fine powder 118 before returning to the input gas stream. Ascan be seen from FIG. 6A, the surroundings to which the melt andeventual powder are exposed are completely controlled.

There are many effective nozzle designs known in the art to producespherical metal powder. Nozzle designs with short gas-to-melt separationdistances produce finer powders. Confined nozzle designs where gas meetsthe molten stream at a short distance just after it leaves theatomization nozzle are preferred for the production of the inventive L1₂aluminum alloy powders disclosed herein. Higher superheat temperaturescause lower melt viscosity and longer cooling times. Both result insmaller spherical particles.

A large number of processing parameters are associated with gasatomization that affect the final product. Examples include meltsuperheat, gas pressure, metal flow rate, gas type, and gas purity. Ingas atomization, the particle size is related to the energy input to themetal. Higher gas pressures, higher superheat temperatures and lowermetal flow rates result in smaller particle sizes. Higher gas pressuresprovide higher gas velocities for a given atomization nozzle design.

To maintain purity, inert gases are used, such as helium, argon, andnitrogen. Helium is preferred for rapid solidification because the highheat transfer coefficient of the gas leads to high quenching rates andhigh supersaturation of alloying elements.

Lower metal flow rates and higher gas flow rates favor production offiner powders. The particle size of gas atomized melts typically has alog normal distribution. An example of spherical L1₂ aluminum alloypowder is shown in the scanning electron micrograph (SEM) of FIG. 7.

Oxygen and hydrogen in the powder can degrade the mechanical propertiesof the final part. It is preferred to limit the oxygen in the L1₂ alloypowder to about 1 ppm to 2000 ppm. Oxygen is intentionally introduced asa component of the helium gas during atomization. A thin oxide coatingon the L1₂ aluminum powder is beneficial for two reasons. First, thecoating prevents agglomeration by contact sintering and secondly, thecoating inhibits the chance of explosion of the powder. A controlledamount of oxygen is important in order to provide good ductility andfracture toughness in the final consolidated material. Hydrogen contentin the powder is controlled by ensuring the dew point of the helium gasis low. A dew point of about minus 50° F. (minus 45.5° C.) to minus 100°F. (minus 73.3° C.) is preferred.

In preparation for final processing, the powder is classified accordingto size by sieving. To prepare the powder for sieving, if the powder haszero percent oxygen content, the powder may be exposed to nitrogen gaswhich passivates the powder surface and prevents agglomeration. Finerpowder sizes result in improved mechanical properties of the endproduct. While minus 325 mesh (about 45 microns) powder can be used,minus 450 mesh (about 30 microns) powder is a preferred size in order toprovide good mechanical properties in the end product. During theatomization process, powder is collected in collection chambers in orderto prevent oxidation of the powder. Collection chambers are used at thebottom of atomization chamber 104 as well as at the bottom of cyclonecollector 116. The powder is transported and stored in the collectionchambers also. Collection chambers are maintained under positivepressure with nitrogen gas which prevents oxidation of the powder.

The process of consolidating the inventive alloy powders into usefulforms is schematically illustrated in FIG. 8. L1₂ aluminum alloy powders210 are first classified according to size by sieving (step 220). Fineparticle sizes are required for optimum mechanical properties in thefinal part. Next, the classified powders are blended (step 230) in orderto maintain microstructural homogeneity in the final part. Blending isnecessary because different atomization batches produce powders withvarying particle size distributions. Other benefits of blending will bediscussed later. Powders may be optionally cryomilled (step 240) tominimize grain size and improve strength. Cryomilling is carried out ina high-energy ball mill under liquid nitrogen, and offers severalbenefits that will be discussed later.

The sieved, blended and (optionally) cryomilled powders are then put ina can (step 250) and vacuum degassed (step 260). Following vacuumdegassing, the can is sealed (step 270) under vacuum and forged (step280) to produce a densified preform. Finally, the preform is Ceraconforged (step 290) to produce a product with improved mechanicalproperties useful for subsequent service as a high temperature L1₂strengthened aluminum alloy. Non-isostatic Ceracon forging will bedescribed later.

Sieving (step 220) is a critical step in consolidation because the finalmechanical properties relate directly to the particle size. Finerparticle size results in finer L1₂ particle dispersion and finer grainsize. Optimum mechanical properties have been observed with −450 mesh(30 micron) powder. Sieving (step 220) also limits the defect size inthe powder. Before sieving, the powder is passivated with nitrogen gasin order to prevent agglomeration. Ultrasonic sieving is preferred forits efficiency.

Blending (step 230) is another critical step in the consolidationprocess because it results in improved uniformity of particle sizedistribution. Gas atomized L1₂ aluminum alloy powder generally exhibitsa bimodal particle size distribution and cross blending of separatepowder batches tends to homogenize the particle size distribution.Blending (step 230) is also necessary when separate metal and/or ceramicpowders are added to the L1₂ base powder to form bimodal and trimodalconsolidated alloy microstructures.

Cryomilling (step 240) can be used to refine the grain size of gasatomized L1₂ aluminum alloy powder as well as the final consolidatedalloy microstructure. Cryomilling is described in U.S. Pat. No.6,902,699, Fritzemeier et al. and in U.S. Pat. No. 7,344,675, Van Daamet al. and are incorporated herein in their entirety by reference.Cryomilling involves high-energy ball milling under liquid nitrogen. Theliquid nitrogen environment prevents oxidation and prevents frictionalheating of the powder and the resulting grain coarsening. During theprocess, the powder particles are repeatedly sheared, fractured and coldwelded which results in a severely deformed microstructure containing ahigh dislocation density that, with continued deformation, evolves intoa cellular structure consisting of extremely small dislocation freegrains separated by high angle grain boundaries with high dislocationdensity. The grain size of the cellular microstructure is typically lessthan 100 nm and the microstructure is considered a nanostructure.

In addition, the nitrogen environment results in the formation ofnitride particles that reside at the grain boundaries and in the graininteriors and resist coarsening at higher temperatures. Stearic acid ispreferably added to the powder charge to prevent excessive agglomerationand to promote fracturing and rewelding of the L1₂ aluminum alloyparticles during milling.

Following sieving (step 220), blending (step 230) and (optionally)cryomilling (Step 240), the powders are transferred to a can (step 250)where the powder is vacuum degassed (step 260) for about 12 hours toover 8 days at elevated temperatures. A temperature range of about 500°F. (260° C.) to about 900° F. (482° C.) is preferred and about 750° F.(399° C.) is more preferred. Dynamic degassing large amounts of powderare preferred to static degassing to expose all of the powder to auniform temperature. Degassing removes the stearic acid lubricant aswell as oxygen and hydrogen from the charge.

Following vacuum degassing (step 260), the vacuum line is crimped andwelded shut. The powder is then consolidated into a dense preform byclosed die forging or by quasi-isostatic forging (step 280).

An exemplary embodiment of this invention is to consolidate a canned L1₂aluminum alloy powder preform into a substantially 100% dense billet bya quasi-isostatic Ceracon-type forging process. The forging processconsists of uniaxially pressing the canned powder preform or solid partperform in a cylindrical press, wherein the preform is surrounded bypressure transmitting medium during the forging. A schematic ofquasi-isostatic forging equipment 300 is shown in FIG. 9. The equipmentconsists of cylindrical die 310 on base 320 with ram 325 inserted in dyecavity 330. Quasi-isostatic forging is carried out at elevatedtemperatures schematically illustrated by heating coils 340.

The steps to Ceracon forge an L1₂ aluminum alloy powder preform areschematically illustrated in FIG. 10. First die 310 is (optionally)preheated (step 410). The preheat temperature is preferably about 40° F.(4.5° C.) to about 70° F. (21° C.) higher than the forging temperatureto compensate for cooling that occurs during the loading process. Next,die 310 is partially filled with pressure transmitting medium (PTM) 360(step 420). To minimize cooling during loading, the PTM can be(optionally) preheated, preferably to the same temperature as the die(step 450). Consolidated powder preform 350 is then inserted in die 310.Powder preform 350 can be (optionally) preheated, preferably to the sametemperature as die 310 and PTM 360 (step 440). The remainder of diecavity 330 is then filled with PTM 360 (step 460). PTM 360 can be(optionally) heated (step 450). Ram 325 is then inserted in die 310(step 470). Pressure 370 is applied to Ram 325 to accomplishquasi-isostatic Ceracon type forging.

One advantage of this process is that, if the preheating steps arefollowed, the run time is short. As a result, deleteriousmicrostructural changes, such as grain and particle coarsening areminimized. Run times can be as short as a few minutes if automatedloading and charging equipment is used. In addition, ductility andfracture toughness of L1₂ based aluminum alloys can be improvedsignificantly due to dynamic bimodal pressure generated duringquasi-isostatic forging that breaks apart continuous powder surfaceoxide that prevent metal-metal bonding and randomly distributes them inthe material.

During quasi-isostatic Ceracon type forging, the billet axially deformsabout 30% and radially deforms about 10%. Strain rates from about 0.1min⁻¹ to about 6 min⁻¹ at forging temperatures from about 400° F. (204°C.) to about 900° F. (482° C.) are preferred. Forging pressures from 50ksi (345 MPa) to 150 ksi (1034 MPa) are preferred.

The pressure transmitting medium can consist of graphite or other carboncontaining powders or ceramic powders or both. By proper fixturing andpreheating the pressure transmitting medium and the work piece,densification during quasi-isostatic forging can be extremely rapid onthe order of minutes.

The unequal actual and radial deformation of the powder is schematicallyillustrated by the dotted line outline of deformed can 352 in FIG. 9.

Quasi-isostatic forging of a 60 to 80 percent dense aluminum-7.5 weightpercent magnesium powder preform by this technique resulted in about a30 percent axial compression and about a 10 percent radial expansion astaught by Meeks III et al. U.S. Pat. No. 7,097,807 and included hereinby reference. The quasi-isostatic forging deformation of L1₂ aluminumalloy powder of this embodiment has beneficial effects on themicrostructure and resulting properties of the consolidated billet. Thenonuniform stress and resulting strain field in the powder duringforging results in extensive shear deformation. The shear deformationdeforms the powder, strips off the surface oxide from the L1₂ alloyparticles and redistributes it throughout the consolidated L1₂ aluminumalloy powder forging as finely divided dispersoids. As a result, thereis increased metal to metal contact during forging and resultingincreased mechanical integrity. The redistributed surface oxideparticles act as additional strengthening agents by resistingdislocation motion by Orowan strengthening.

In other embodiments, Ceracon type forging can be followed by hotisostatic pressing (HIP), forging, rolling and other deformationprocessing techniques.

To summarize, Ceracon type quasi-isostatic forging can be low cost andefficient. The pressure transmitting medium, as well as the vacuumsealed preform can be preheated prior to introduction of the performinto the forging chamber to minimize runtime. In addition, forging canbe accomplished at high strain rates resulting in forging runs lastingless than a few minutes. The short run time results in an economicalprocess that also further inhibits grain and L1₂ particle growth in thebillet during forging due to the limited time at temperature. Thepressure transmitting medium can be reused and loading and unloading thepreforms and resulting forgings can be an automated procedure.

Although the present invention has been described with reference topreferred embodiments, workers skilled in the art will recognize thatchanges may be made in form and detail without departing from the spiritand scope of the invention.

1. A method for producing high strength aluminum alloy consolidatedbillets containing L1₂ dispersoids, comprising the steps of: forming analuminum alloy powder containing L1₂ dispersoid forming elements,wherein the L1₂ dispersoid forming elements form Al₃X dispersoidswherein X is at least one first element selected from the groupcomprising: about 0.1 to about 4.0 weight percent scandium, about 0.1 toabout 20.0 weight percent erbium, about 0.1 to about 15.0 weight percentthulium, about 0.1 to about 25.0 weight percent ytterbium, and about 0.1to about 25.0 weight percent lutetium; at least one second elementselected from the group comprising: about 0.1 to about 20.0 weightpercent gadolinium, about 0.1 to about 20.0 weight percent yttrium,about 0.05 to about 4.0 weight percent zirconium, about 0.05 to about10.0 weight percent titanium, about 0.05 to about 10.0 weight percenthafnium, and about 0.05 to about 5.0 weight percent niobium; and thebalance substantially aluminum; placing the powder in a container;vacuum degassing the container at about 500° F. (260° C.) to about 900°F. (482° C.) for about 12 hours to about 8 days; sealing the container;creating a preform by compressing the container by closed die forging orby quasi-isostatic forging to consolidate the powder; encompassing thealuminum alloy powder preform with a flowable pressure transmittingmedium and heating the encompassed alloy powder; and uniaxiallycompressing the medium to thereby consolidate the aluminum powder; andremoving the consolidated powder billet.
 2. The method of claim 1,wherein the forging temperature is from about 400° F. (204° C.) to about900° F. (482° C.).
 3. The method of claim 1, wherein the axial strainrate during compressing the medium is from about 0.1 min⁻¹ to 6 min⁻¹.4. The method of claim 1, wherein compressing the medium is at apressure from about 50 Ksi (345 MPa) to about 150 Ksi (1034 MPa).
 5. Themethod of claim 1, wherein the preform is vacuum sealed in an aluminumjacket.
 6. The method of claim 1 wherein the pressure transmittingmedium comprises graphite or other carbon containing powders or ceramicpowders or both.
 7. The method of claim 1 wherein the aluminum alloypowder contains at least one third element selected from the groupconsisting of silicon, magnesium, lithium, copper, zinc and nickel. 8.The method of claim 1 wherein the third element comprises at least oneof about 4 to about 25 weight percent silicon, about 1 to about 8 weightpercent magnesium, about 0.5 to about 3 weight percent lithium, about0.2 to about 6.5 weight percent copper, about 3 to about 12 weightpercent zinc, and about 1 to about 12 weight percent nickel.
 9. A highstrength aluminum alloy consolidated billet containing L1₂ dispersoidsin an aluminum alloy matrix wherein: the L1₂ dispersoid forming elementsform Al₃X dispersoids wherein X is at least one first element selectedfrom the group comprising: about 0.1 to about 4.0 weight percentscandium, about 0.1 to about 20.0 weight percent erbium, about 0.1 toabout 15.0 weight percent thulium, about 0.1 to about 25.0 weightpercent ytterbium, and about 0.1 to about 25.0 weight percent lutetium;at least one second element selected from the group comprising: about0.1 to about 20.0 weight percent gadolinium, about 0.1 to about 20.0weight percent yttrium, about 0.05 to about 4.0 weight percentzirconium, about 0.05 to about 10.0 weight percent titanium, about 0.05to about 10.0 weight percent hafnium, and about 0.05 to about 5.0 weightpercent niobium; and the balance substantially aluminum, wherein: thealuminum alloy consolidated billet containing L1₂ dispersoids is formedby the steps comprising: placing the powder in a container; vacuumdegassing the container at about 500° F. (260° C.) to about 900° F.(482° C.) for about 12 hours to about 8 days; sealing the container;creating a preform by compressing the container by closed die forging orby quasi-isostatic forging to consolidate the powder; encompassing thealuminum alloy powder preform with a flowable pressure transmittingmedium and heating the encompassed alloy powder; and encompassing thealuminum alloy powder preform with a flowable pressure transmittingmedium and heating the encompassed alloy powder and medium; anduniaxially compressing the medium to thereby consolidate the aluminumpowder; and removing the consolidated powder billet.
 10. The alloy ofclaim 9, wherein the aluminum alloy powder contains at least one thirdelement selected from the group consisting of silicon, magnesium,lithium, copper, zinc, and nickel.
 11. The alloy of claim 10, whereinthe third element comprises at least one of about 4 to about 25 weightpercent silicon, about 1 to about 8 weight percent magnesium, about 0.5to about 3 weight percent lithium, about 0.2 to about 6.5 weight percentcopper, about 3 to about 12 weight percent zinc, and about 1 to about 12weight percent nickel.
 12. The alloy of claim 9, wherein the forgingtemperature is from about 400° F. (204° C.) to about 900° F. (482° C.).13. The alloy of claim 9, wherein the axial strain rate duringcompressing the medium is from about 0.1 min⁻¹ to 6 min⁻¹.
 14. The alloyof claim 9, wherein compressing the medium is at a pressure from about50 Ksi (345 MPa) to about 150 Ksi (1034 MPa).
 15. The alloy of claim 9,wherein the pressure transmitting medium comprises graphite or othercarbon containing powders or ceramic powders or both.
 16. A highstrength aluminum alloy consolidated billet containing L1₂ dispersoidsin an aluminum alloy matrix wherein: the L1₂ dispersoid forming elementsform Al₃X dispersoids wherein X is at least one first element selectedfrom the group comprising: about 0.1 to about; 4.0 weight percentscandium, about 0.1 to about 20.0 weight percent erbium, about 0.1 toabout 15.0 weight percent thulium, about 0.1 to about 25.0 weightpercent ytterbium, and about 0.1 to about 25.0 weight percent lutetium;at least one second element selected from the group comprising: about0.1 to about 20.0 weight percent gadolinium, about 0.1 to about 20.0weight percent yttrium, about 0.05 to about 4.0 weight percentzirconium, about 0.05 to about 10.0 weight percent titanium, about 0.05to about 10.0 weight percent hafnium, and about 0.05 to about 5.0 weightpercent niobium; and the balance substantially aluminum, wherein: thealuminum alloy consolidated billet containing L1₂ dispersoids is formedby the steps comprising: placing the powder in a container; vacuumdegassing the container at about 500° F. (260° C.) to about 900° F.(482° C.) for about 12 hours to about 8 days; sealing the container;creating a preform by compressing the container by closed die forging orby quasi-isostatic forging to consolidate the powder; encompassing thealuminum alloy powder preform with a flowable pressure transmittingmedium and heating the encompassed alloy powder and medium; anduniaxially compressing the medium at an axial strain rate of from about0.1 min⁻¹ to about 6 min⁻¹ at a pressure from about 50 KSi (345 MPa) toabout 150 KSi (1034 MPa) to thereby consolidate the aluminum powder; andremoving the consolidated powder in a billet.
 17. The high strengthaluminum alloy consolidated billet of claim 16 wherein the aluminumalloy powder contains at least one third element selected from the groupconsisting of silicon, magnesium, lithium, copper, zinc, and nickel. 18.The high strength aluminum alloy consolidated billet of claim 16 whereinthe third element comprises at least one of about 4 to about 25 weightpercent silicon, about 1 to about 8 weight percent magnesium, about 0.5to about 3 weight percent lithium, about 0.2 to about 6.5 weight percentcopper, about 3 to about 12 weight percent zinc, and about 1 to about 12weight percent nickel.
 19. The high strength aluminum alloy consolidatedbillet of claim 16 wherein the forging temperature is from about 400° F.(204° C.) to about 900° F. (482° C.).
 20. The high strength aluminumalloy consolidated billet of claim 16 wherein the pressure transmittingmedium comprises graphite or other carbon containing powders or ceramicpowders or both.